structure-property relationships directing transport and ... · 1 structure-property relationships...
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Structure-Property Relationships Directing
Transport and Charge Separation in Isoindigo
Polymers
Caroline Grand,1 Sujin Baek,2 Tzung-Han Lai,2 Nabankur Deb,3 Wojciech Zajaczkowski,4
Romain Stalder,5 Klaus Müllen,4 Wojciech Pisula,4,6 David G. Bucknall,3 Franky So,2† John R.
Reynolds1*
1. School of Chemistry and Biochemistry, School of Materials Science and Engineering, Center
for Organic Photonics and Electronics, Georgia Tech Polymer Network, Georgia Institute of
Technology, Atlanta, Georgia 30332-0400, United States.
2. Department of Materials Science and Engineering, University of Florida, Gainesville, Florida
32611-7200, United States.
3. School of Materials Science and Engineering, Center for Organic Photonics and Electronics,
Georgia Institute of Technology, Atlanta, Georgia 30332-0400, United States.
4. Max Planck Institute for Polymer Research, Ackermannweg 10, 55128 Mainz, Germany.
5. The George and Josephine Butler Polymer Research Laboratory, Department of Chemistry,
University of Florida, Gainesville, Florida 32611-7200, United States.
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6. Department of Molecular Physics, Faculty of Chemistry, Lodz University of Technology,
Zeromskiego 116, 90-924 Lodz, Poland.
KEYWORDS. Polymer Solar Cells, Isoindigo Polymers, Bulk Heterojunctions
ABSTRACT. Since being introduced to the open literature in 2010, the isoindigo heterocycle has
been extensively studied as a novel electron deficient building block for organic electronic
materials in conjugated polymers, discrete length oligomers, and molecular systems, particularly
targeting high charge mobility values and ambipolar transport in organic field effect transistors,
along with high power conversion efficiencies in organic photovoltaic devices. This article
introduces results obtained on copolymers of isoindigo with thiophene and alkylated
terthiophenes to highlight fundamental characteristics in isoindigo-based polymers and the
resulting organic field-effect transistors and photovoltaic devices. By comparing and contrasting
the optoelectronic properties, thin film morphology, organic field-effect transistor (OFET)
mobilities, and organic photovoltaic (OPV) performance to previously reported polymers,
structure-processing-property relationships were uncovered. In particular, isoindigo-containing
polymers with more rigid backbones and higher coherence lengths in thin-films lead to increased
charge mobility in OFET devices. In OPV devices, efficiencies over 6% can be obtained by
balancing high ionization potentials typically dictating the open-circuit voltage, and the charge
transfer state energy and blend morphology impacting short-circuit currents. Furthermore, the
impact of polymer structure on solubility and on phase separation in blends with PC71M is
discussed, with isoindigo-based polymers exhibiting lower solubility possibly leading to more
fiber-like morphologies due to both kinetic and thermodynamic effects, either stemming from
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polymer dissolution in the casting solvent or from polymer self-assembly during film formation.
This fiber-like polymer morphology remains unaffected by the presence of processing additives,
such as 1,8-diiodooctane. These structure-property relationships developed for isoindigo-based
polymers can also be discussed in the broader context of diketopyrrolopyrrole (DPP) and
thienoisoindigo (TiI) as electron-deficient moieties that can also be doubly substituted on their
amide functionality.
INTRODUCTION. π-Conjugated polymers have been successfully used in solution processed
organic field-effect transistors (OFETs) and organic photovoltaic (OPV) devices, with FET
measured hole mobilities1, 2 on the order of 10 cm2 V-1 s-1 and electron mobilities3, 4 ranging from
0.1 to 1.0 cm2 V-1 s-1,5 while solar cells have attained AM 1.5 power conversion efficiencies
(PCEs) in excess of 9.2%,6-11 and 10.5% in tandem cells.12 In examining this field of materials,
similar polymer structures do not translate from one application to another, but rather need to be
designed to fulfill application specific requirements. In materials designed for OFETs for
instance, there is a need to combine (i) solution processability, (ii) high charge carrier transport,
and (iii) ambient air stability. On the other hand, materials for OPVs may not require record-high
carrier mobilities to perform well, but they also need (iv) broad optical absorption profiles and
high extinction coefficients, as well as (v) appropriate energy levels for charge separation.
Moreover, the solid state order in OFETs,13-16 and phase separation in OPVs are crucial
components towards high performing devices.17
Donor-acceptor (D-A) polymers have been developed using electron accepting units such as
1,3,2-benzothiadiazole (BTD), thienopyrrolodione (TPD) or diketopyrrolopyrrole (DPP) to yield
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tunable optoelectronic properties. Recently, the electron deficient isoindigo (iI) unit has shown
promise as an electron acceptor in both OFETs and OPV materials, thanks to its high yielding
synthetic access, two amide functionalities for dual substitution to ensure polymer solubility,
large absorption coefficients and to the suitable electron affinity of the isoindigo polymers to be
used in fullerene-based OPV devices.18 The iI moiety was first reported in a CIBA patent,19 and
subsequently introduced into the open literature by Reynolds et al.20 Since then iI-based
polymers have demonstrated FET hole mobilities21, 22 up to 3 cm2 V-1 s-1 thanks to side-chain
engineering using longer spacers to direct branching points away from the polymer backbone, as
was also reported for other D-A polymers.23 In OPV devices, efficiencies24-27 of 7.0% were
reached using polymer structures based on thienylene-phenylene linkages by tuning the energy
of the frontier molecular orbitals to achieve a balance between a deep electron affinity for
efficient light absorption, charge transfer and charge separation for high short-circuit current
(Jsc), and a deep ionization potential for a high open-circuit voltage (Voc).28 Isoindigo is an
attractive unit for organic electronics as it possesses two amide groups that can be functionalized,
similar to DPP, but in contrast with BTD derivatives, which do not have similar groups for
substitution, and TPD units, which only have one handle for substitution. The possibility of
functionalizing the iI core with solubilizing alkyl, oligoether or siloxane-terminated chains
provides the opportunity to tune optoelectronic and morphological properties, as is the case for
DPP or thienoisoindigo (TiI) units.23
Based on previous literature on iI-based polymers, six fundamental principles are made evident
in these systems with the goal of understanding the correlation between polymer structure, solid-
state morphology, and device properties. While this paper, and the analysis listed below, is
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directed to the study of iI-based polymers specifically, similar concepts can be drawn to other bi-
substituted acceptors like DPP and TiI:
- The electron affinity remains around 3.9 eV, as the electronic density in the LUMO is
located on the isoindigo core, while the ionization potential can be tuned by the donor
moiety.29
- Control of the ionization potential leads to tunable open-circuit voltages in OPV devices,
ranging from 0.6 V to 0.9 V depending on the donor selected,18, 30 correlating well with the
effective band gap, i.e. excitation from the ground state of the donor to the excited state of
the acceptor.
- The packing and orientation of the polymer chains can be influenced both by strong
interchain interactions made possible by the rigidity of the isoindigo core compared to
phenylene-flanked DPP, and by the side chain choice.22, 31
- Isoindigo-based polymers can lead to ambipolar transport,32 and high transport
characteristics (up to 3.6 cm2 V-1 s-1 hole mobility,21 and 0.7 cm2 V-1 s-1 electron mobility in
chlorinated isoindigo polymers33). An increase in mobility values is not necessarily
correlated to a decrease in π-stacking distances, and identical characteristic distances can
lead to different mobilities if the morphology is different.
- Polymer solubility can be tuned through the side-chain density along the polymer
backbone as well as through the nature of the side-chains.34, 35 Furthermore, phase
separation between polymer and fullerene in the solid-state is dictated by polymer
solubility in the processing solvent,36 and phase separation is not as drastically modified by
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the use of processing additives when polymers exhibit low solubility (<2 mg/mL in ortho-
dichlorobenzene).
- In bulk heterojunctions with PCBM derivatives, polymer-rich domain sizes on the order
of tens of nanometers surrounded by mixed regions of polymer and fullerene can be
achieved through structural design and processing control.37 When coupled to low
resistance in OPV devices this leads to high short-circuit currents and fill factors.
While these principles have been applied to different structures very rarely are they considered
together. To illustrate these concepts, and in particular highlight the donor-acceptor effects on
frontier energy levels and the impact of side-chains on polymer packing and morphology, we
designed a family of three oligothiophene-isoindigo polymers with varying ratios of electron
deficient iI units along the backbone, and variations in the solubilizing side-chain as shown in
Chart 1. The energy levels of the frontier molecular orbitals of an all-thiophene backbone are
modulated by the content of isoindigo units along the conjugated backbone to shift the
absorption to longer wavelengths going from P[T-iI(HD)] to P[T3(C6)-iI(HD)]. Furthermore,
while keeping the conjugated repeat unit constant and following on the work of Bao and co-
workers,22, 38 we designed P[T3(C6)-iI(SiO)] to investigate the effect of the substitution pattern
on solid state order, transport properties and phase separation.
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Chart 1. Structures of P[T-iI(HD)], P[T3(C6)-iI(HD)], and P[T3(C6)-iI(SiO)] (red boxes, with
new results discussed in this article), compared to previously published polymers. The color
scheme highlights structures with the same conjugated repeat unit: P[T-iI] – black, P[T2-iI] –
orange; P[DTS-iI] – brown; P[T3-iI] – blue; P[T6-iI] – green.
In this contribution, three new isoindigo polymers are studied in OFET and OPV devices, and
their behavior is compared and contrasted with previously published isoindigo-based polymers
described in Chart 1. These comparative results are used to elucidate a set of guidelines for
molecular design principles for isoindigo containing polymers in organic electronic applications.
The structures discussed investigate the impact of electron-richness along the P[Tn-iI] backbone
on optoelectronic properties, as the electron-rich co-monomer is extended from P[T-iI] (black),
to P[T2-iI] (orange), P[DTS-iI] (brown), P[T3-iI] (blue) and P[T6-iI] (green). With the repeat
unit kept constant, the side-chains on P[Tn(R)-iI(R)] can also be modified to induce changes in
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the polymer packing and optoelectronic properties. This study makes detailed comparison of
optoelectronic structure-property relationships to work published on P[T-iI(OD)] from Zhang et
al.,39 P[T-iI(HD)], P[T3(C8)-iI(HD)], and P[T6(C8)-iI(HD)] from Ma, Wang et al.25, 28, 37, 40, 41
and Ho et al.,26 P[T2-iI(OD)] and P[T2-iI(SiO)] from Mei, Kim et al.,22, 42, P[T-iI(OD)], P[T2-
iI(OD)] and P[T2-iI(4-DTd)] from Lei et al.,21, 43 and P[DTS(EH)-iI(EH)] from Stalder et al.44
Grazing incidence wide angle X-ray scattering (GIWAXS) is used to emphasize the role of
microscopic order and backbone orientation on charge transport in OFETs. Density functional
theory (DFT) calculations, along with electrochemical results and charge-modulated
electroabsorption spectroscopy (CMEAS) measurements, were used to understand the charge
separation process and ultimately the open-circuit voltage in OPV devices. Furthermore, AFM
and neutron reflectivity were used to study the lateral and vertical phase separation in the active
layer blends with fullerene derivatives to explain variation in Jsc and fill factor (FF) in OPV
devices.
RESULTS AND DISCUSSION
Polymer synthesis. Isoindigo-based polymers have typically been synthesized using Suzuki and
Stille polycondensations. The synthesis of P[T-iI(HD)]-4 was previously reported,32 and
P[T3(C6)-iI(HD)]-2 and P[T3(C6)-iI(SiO)] were synthesized by Stille polymerization as
described in Scheme S1. We aimed to increase ordering of P[T3-iI] polymers by using n-hexyl
chains on the terthiophene unit rather than the n-octyl chains used previously40 to investigate the
impact of alkyl side-chain length on polymer packing for OFET applications, along with
examining the effect on phase separation with fullerenes in bulk heterojunction (BHJ) OPV
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devices. By moving the branching point of the 2-hexyldecyl side chain further away from the
backbone using siloxane side chains, as demonstrated by Mei et al.,45 the impact of polymer
orientation on its interaction with fullerenes can be determined. The number average molecular
weight for these three polymers was estimated by gel permeation chromatography (GPC) in
tetrahydrofuran (THF) and correlates with a degree of polymerization (DP) of 34, 16, and 25 (i.e.
102, 80, and 125 aromatic rings) for P[T-iI(HD)]-4, P[T3(C6)-iI(HD)]-2 and P[T3(C6)-iI(SiO)],
respectively (see Table 1: entries 6, 14 and 19, and SI). The three polymers also have comparable
indices of polydispersity in THF (2.1, 3.7, and 2.2, respectively). Important for comparison,
these macromolecules are similar in size and have reached a point where saturation of their
optoelectronic properties is ensured. In particular, the use of branched side chains on the
isoindigo unit for these polymers ensures solubility during polymerization, and prevents
precipitation at lower molecular weights during polymerization. Furthermore, all three polymers
structures were synthesized to have thienyl end-groups; however, the presence of controlled end-
cappers was not confirmed experimentally.
Theoretical Electronic and Structural Analysis. Estrada et al.29 have previously demonstrated
good correlation between HOMO and LUMO energy levels calculated by density functional
theory (DFT) at the B3LYP/6-31G(d) level and electrochemically determined ionization
potential and electron affinity. As such, DFT calculations at the B3LYP/6-31G(d) level were
carried out on model structures of P[T-iI] and P[T3-iI] (where side chains have been replaced by
methyl groups) to investigate shifts in the energy of the frontier molecular orbitals as a function
of the electron-rich oligothiophene co-monomers. As expected, the less electron-rich T-iI motif
has a deeper calculated HOMO level at -4.72 eV compared to -4.97 eV for T3-iI. This 0.25 eV
difference in HOMO level energy is expected to be reflected in the open circuit voltage (Voc) of
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OPV devices with polymers based on a T-iI repeat unit having an increased Voc compared to T3-
iI polymers. In both T-iI and T3-iI, the HOMO is highly delocalized along the polymer chain,
whereas the LUMO is mostly localized on the isoindigo unit as demonstrated by the T3-iI dimer
in Figure 1. The optimized ground state geometry in these model compounds shows a break in
conjugation in the LUMO of P[T3-iI], with electron density located over a single isoindigo unit
and extending slightly over the neighboring thiophene rings. In P[T-iI] systems, the electron
density in the LUMO is spread continuously over two isoindigo cores bridged by a thiophene
ring as seen in Figure S1, which would modify the intrachain transport of electrons and possibly
their interchain hopping characteristics.46, 47 Similar DFT results for T-iI oligomers have been
obtained by De Angelis et al.,48 where the calculated HOMO and LUMO levels (in vacuum) are
within 30 meV of those found in this work, despite a different conformation used for modeling
of T-iI oligomers.
Figure 1. Isodensity surfaces (0.04 e/bohr3) of the frontier orbitals a T3-iI dimer calculated at the
B3LYP/6-31G(d) level: (a) LUMO and (b) HOMO (adapted from 49).
Electrochemically Determined Energy Levels, and Optical Properties. To verify the trends
observed in the energy levels of the frontier molecular orbitals, electrochemical measurements of
the oxidation and reduction potentials were conducted against a ferrocenium/ferrocene (Fc+/Fc)
standard to estimate the ionization potential and electron affinity of the polymers. Based on
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measurements of the formal potential for Fc+/Fc at 0.4 V versus a saturated calomel electrode
(SCE) in acetonitrile,50 and of -4.5 eV being equivalent to 0.0 V versus a normal hydrogen
electrode (NHE)51-53 (with SCE at +0.2 V versus NHE ). The formal potential of Fc+/Fc is taken
to be -5.1 eV versus vacuum, which was highlighted by Cardona et al.54 Differential pulse
voltammetry (DPV) measurements on polymer thin films on platinum electrodes in 0.1 M
TBAPF6 in acetonitrile estimated the electron affinity (EA) of the three polymers to be between
3.9 and 3.94 eV. The values match previous reports of the EA of donor-isoindigo-donor
structures being around 3.7 eV and 3.8 eV,18 and are in agreement with the EA for other iI
copolymers listed in Table 1. The ionization potential (IP) of these structures was estimated from
the onset of oxidation. The IP was calculated to be 5.59 eV, 5.57 eV and 5.43 eV for P[T-
iI(HD)]-4, P[T3(C6)-iI(HD)]-2 and P[T3(C6)-iI(SiO)], respectively (entries 6, 14 and 20 in
Table 1). Previously reported values for the HOMO energy level demonstrate that the measured
IP varies with the technique used, with IP of P[T2-iI(OD)] (entry 9) measured at 5.5 eV using
UV-PES and 6.0 eV using cyclic voltammetry with ferrocene oxidation at -5.1 eV against
vacuum.54 In order to support the IP estimated from DPV measurements, UV-PES was also
conducted on thin-films of P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2. The IP was found to be
around 5.2 eV and 5.3 eV in films of P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2 respectively, again
highlighting a 0.4-0.3 eV difference between the IP estimated by DPV relative to the IP
measured by UV-PES. Furthermore, the electron affinity of these polymers is estimated to be
between 3.8 eV and 4.1 eV as expected based on previous DFT calculations, where the
calculated electron density in the LUMO was primarily localized on the iI unit within model
compounds.29 The IP fluctuated over a broader energy range (5.3 eV to 6.1 eV) based on the
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electronic nature of the donor moiety, where iI-based polymers with more electron-rich
thiophene units tend to exhibit lower IP’s.
As observed in Figure 2, the absorption spectra of all three isoindigo polymers show a broad
coverage of the visible region, from 400 to 800 nm with absorption maxima around 650 nm.
There are only minor differences observed between the solution and solid-state spectra, where
often times in conjugated polymers a substantial red-shift is observed in the solid state spectra
compared to the spectra in solution. We attribute this to minimal changes in polymer
conformation upon solidification, as both planarization of the polymer chains and interchain
interactions in the solid state lead to a red-shift of the low energy absorption peak or to the
appearance of a low energy shoulder in the UV-vis spectrum as discussed by Kim and Swager.55
Broader UV-vis spectra with a low energy aggregation shoulder are observed for films of
P[T3(C6)-iI(HD)]-2 and P[T3(C6)-iI(SiO)] compared to solution, suggesting an increase in the
intermolecular interactions between neighboring polymer chains in the solid state compared to
solution state. This low energy shoulder is not observed in the case of P[T-iI(HD)]-4 compared
to solution, suggesting a more amorphous arrangement of the polymer backbones in thin films.
This decrease in thin film order for P[T-iI(HD)]-4 is hypothesized to be linked to the higher
density of branched alkyl chains along the backbone, decreasing the extent of interchain
interactions. This discussion based on UV-vis-NIR results is also supported by 2D-NMR
spectroscopy and X-ray scattering data (vide intra). Making comparisons of the many iI-based
polymers reported, all of the entries in Table 1 have an onset of absorption between 1.50 eV and
1.62 eV, likely due to the similar electronic nature of the thiophene-based donors along the
backbone, also described by Stalder et al.18 Figure 2 also highlights that the solution and solid-
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state absorption of these oligothiophene-co-iI polymers exhibit two main absorption bands,
which span the visible region of the spectrum to yield blue-green colored polymers.
Table 1. Size exclusion chromatography (SEC), Redox, and Optical Properties of the Isoindigo
Copolymers
SEC UV-vis-NIR Electrochemistry
Entry Polymer Mn
(kDa) PDI
Emax
d
(eV)
solution
Emax
d
(eV)
film
Egap
opt
(eV)
IPe
(eV)
EAe
(eV)
Egap
(eV) Ref
1 P[T-iI(OD)]-1 20 a 2.0 1.76 1.76 1.55 6.1 4.10 2.00 56
2 P[T-iI(OD)]-2 17a 2.1 1.81 1.79 1.58 5.88
39
3 P[T-iI(HD)]-1 86b 2.2 1.78 1.78 1.6 5.82 3.85 1.83 41
4 P[T-iI(HD)]-2 48b 2.7
1.82
5.72 4.00
28
5 P[T-iI(HD)]-3 60b 2.2
1.55 5.75 3.92 1.83 25
6 P[T-iI(HD)]-4 26 a 2.1 1.79 1.79 1.62 5.59/5.21
f 3.90 1.69 32
7 P[T2-iI(OD)]-1 88 a 2.1 1.78 1.78 1.55 6.0 4.0 2.00 56
8 P[T2-iI(OD)]-2 135 a 5.5 1.77 1.77 1.62 5.38
f
22
9 P[T2-iI(OD)]-3 20b 2.0 1.77 1.77 1.60 6.00/5.54
f 4.00 2.00 21
10 P[T2-iI(4-DTd)] 39b 3.2 1.73 1.72 1.58 5.82/5.33
f 4.04 1.78 21
11 P[T2-iI(SiO)] 138 a 3.3 1.70 1.74 1.61 5.20
f
22
12 P[DTS(EH)-iI(EH)] 36 a 2.8 1.71 1.72 1.54 5.55 3.95 1.60 44
13 P[T3(C6)-iI(HD)]-1 24 a 1.9 1.92 1.88 1.58 5.39 3.89 1.50 32
14 P[T3(C6)-iI(HD)]-2 18
a 3.7
1.88 1.90 1.58 5.57/5.29f 3.93 1.64
103c
1.6
15 P[T3(C8)-iI(HD)]-1 43b 3.1 1.91 1.92 1.50 5.79 3.80 1.99 40
16 P[T3(C8)-iI(HD)]-2 73b 2.9
1.92
5.58 3.95
28
17 P[T3(C8)-iI(HD)]-3 100b 3.3
1.50 5.66 3.94 1.72 25
18 P[T3(C8)-iI(HD)]-4 48b 2.4 1.92 1.92 1.59 5.49 3.90 1.59 26
19 P[T3(C6)-iI(SiO)] 32a 2.2 1.89 1.88 1.55 5.43 3.94 1.46
14
49c
5.4
20 P[T6(C8)-iI(HD)]-1 67b 2.6
1.52 5.44 3.91 1.53 25
21 P[T6(C8)-iI(HD)]-2 35b 1.8 2.00 1.96 1.57 5.37 3.80 1.57 26
Results from this work are highlighted in red. SEC performed in: aTHF, b1,2,4-trichlorobenzene, coCB, all
versus polystyrene standards, dAbsorption maxima in solution and thin film, eionization potential and
electron affinity estimated from CV or DPV homogenized using EIP(EA) = 5.1 + Eonsetox(red) vs Fc+/Fc ;
fHOMO determined by UV-PES. Values in italic obtained from spectra.
Reflecting on the spectroscopic properties of the full set of iI-based polymers in Table 1, it is
seen that isoindigo P[Tn-iI] copolymers show an aggregation shoulder in thin films for n>3
although broadening of the absorption spectra is seen for all the isoindigo polymers listed. Even
in P[T2-iI(SiO)] and P[T2-iI(4-DTd)] films, where the branching point on the side chain is
moved away from the backbone, no aggregation shoulder is observed in the thin film absorption
spectra. As such, based on the UV-vis spectra alone, it would seem that polymers with a lower
density of branched side chains exhibit stronger interchain interactions, to be correlated to 2D-
NMR and X-ray scattering results.
We recently reported the use of solid-state 2D-NMR spectroscopy to probe the correlation
between branched side chain density and interchain interactions.32 The 2D 13C{1H} FSLG-
HETCOR NMR correlation spectra of powder samples of P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-
1 were compared, effectively allowing us to visualize which protons of the branched side chains
contacted strongly with which carbon of the neighboring polymer backbones. From the 2D solid-
state NMR data alone, the higher branched side chain density for P[T-iI(HD)]-4 correlated with
greater overlap of neighboring isoindigo units in the bulk compared to P[T3(C6)-iI(HD)]-1, for
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which the 2D NMR spectra suggested stronger intermixing of the isoindigo and terthiophene
portions of the polymer backbone. Given the steric hindrance arising from side chain
functionalization in conjugated polymers, this is consistent with a lower density of branched side
chain affording greater interchain interaction, as seen from the UV-vis data in Figure 2.
300 400 500 600 700 800 900
Film
No
rma
lize
d a
bso
rba
nce
Wavelength (nm)
P[T-iI(HD)]-4
P[T3(C6)-iI(HD)]-2
P[T3(C6)-iI(SiO)]Solution
Figure 2. Normalized UV-vis spectra of P[T-iI(HD)]-4 (black), P[T3(C6)-iI(HD)]-2 (red), and
P[T3(C6)-iI(SiO)] (blue) in chloroform solution (top) and thin film (bottom).
Pristine Polymer Thin Film GIWAXS Characterization. To investigate polymer packing and
the interchain interactions mentioned above, grazing incidence wide angle X-ray scattering
(GIWAXS) measurements were performed on thin films prepared by spin-coating from ortho-
dichlorobenzene (oDCB) solutions onto silicon wafers with a layer of silicon oxide at the
surface. The samples were measured both as-cast and dried at room temperature (GIWAXS
measurements made approximately 7 days after preparation) and after annealing at 200 °C (see
Figure 3). Indexing of the scattering peaks was related to the polymer packing structure, with
(100) planes correlating to lamellar stacking and (010) planes to π-stacking. The interplane
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distances were then calculated using the Bragg condition: 𝑑 = 2𝜋𝑄ℎ𝑘𝑙
⁄ .57 After integration over
a quadrant of the detector, line shape analysis was conducted using Gaussian functions to fit the
scattering peaks – the peak positions are attributed to the (hkl) scattering planes and the full
widths at half maximum (FWHM) give indications on the coherence length (CL) calculated by
Scherrer analysis.
Pristine polymer thin films exhibit varying π-stacking distances ranging from 3.6 Å for
P[T3(C6)-iI(SiO)] to 3.8 Å for P[T-iI(HD)]-4, and P[T3(C6)-iI(HD)]-2 has an ordered structure
with a characteristic distance of 4.1 Å that cannot be attributed to π-stacking (Figure 3a-c). This
corroborates what is observed from thin film absorbance measurements with P[T3(C6)-iI(SiO)]
showing the largest red shift among the three polymers. The lack of clear π-stacking scattering
and the larger characteristic distance for P[T3(C6)-iI(HD)]-2 can be explained by increased
rotational freedom in the terthiophene unit containing alkyl chains, compared to a single
thiophene unit in P[T-iI(HD)]-4. The lamellar distance determined from the position of the (100)
peak is similar in films of P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2, around 18 Å, but increases to
25 Å in films of P[T3(C6)-iI(SiO)], possibly due to the longer Si-O-Si distance compared to C-
C-C distance (3.1 Å vs. 2.5 Å, respectively) leading to an overall side chain maximum length of
around 11 Å and 13 Å in the case of 2-hexyldecyl and siloxane side chains respectively. These
values for π-stacking and lamellar distances are close to previous results obtained on
oligothiophene-isoindigo polymers.25 Prior to thermal annealing, the polymer films exhibit
different backbone orientations relative to the substrate. P[T-iI(HD)]-4 and P[T3(C6)-iI(SiO)]
both exhibit (100) lamellar peaks along the Qxy axis (in-plane) and consequently a (010) π-
stacking peak along the Qz axis (polymer backbone mainly face-on relative to the substrate). By
comparison, P[T3(C6)-iI(HD)]-2 has its (100) peak along the Qz axis, and a scattering peak at a
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30° angle from the Qxy axis indicating the presence of scattering layers predominantly tilted
towards, and not edge-on to, the substrate.
Figure 3. GIWAXS patterns of pristine (top) and thermally annealed (bottom) P[T-iI(HD)]-4
(a,d), P[T3(C6)-iI(HD)]-2 (b,e) and P[T3(C6)-iI(SiO)] (c,f) thin films prepared from oDCB.
Figures a-e use the same 5,000 intensity scale bar; for optimal data visualization. Figure f uses
the 10,000 intensity scale bar to visualize lower intensity peaks before annealing, without
saturating the scattering intensity after annealing.
The difference in orientation is made even more visible after thermal annealing at 200°C for 5
minutes as shown in Figure 3d-f. In particular in the case of P[T3(C6)-iI(HD)]-2 films, the
scattering peak at 30° off from the Qxy axis (Figure 3b, arrow) disappears, and two new peaks
18
appear along the Qxy axis after thermal annealing (Figure 3e, arrow). The outermost peak is
attributed to π-stacking of 3.7 Å indicating that the polymer backbones become primarily edge-
on with regards to the substrate after thermal annealing (Table 2, entry 11). The second peak
along the Qxy axis corresponds to a spacing of 4.2 Å, possibly resulting from a rearrangement of
the scattering of scattering peak at 30° off from the Qxy axis. Furthermore, the π-stacking
distance remains at 3.8 Å and 3.6 Å in P[T-iI(HD)]-4 and P[T3(C6)-iI(SiO)] respectively (Table
2: entries 3 and 13 respectively). Thermal annealing also enables a higher degree of order to be
reached in these systems as shown by the higher order scattering peaks in P[T-iI(HD)]-4 and
P[T3(C6)-iI(HD)]-2, and the increase in scattering intensity in films of P[T3(C6)-iI(SiO)]
(Figures 3f uses a different intensity scale than 3e to prevent saturation of the scattering intensity
in the figure). In order to quantify order in the π-stacking direction, Scherrer analysis was used to
calculate CL based on the (010) peak. Interestingly, after thermal annealing the CL of the three
polymers is comparable (around 9-10 nm), corresponding to approximately 20-30 polymer
backbones involved in π-stacks. These numbers are on par with previously reported CL for (010)
peaks for P[T2-iI(OD)]-2 (entry 5), and slightly lower than the length of P[T2-iI(SiO)] (entry 8),
which includes 38 backbones within a π-stack, and higher than P[T3(C8)-iI(HD)]-3 (entry 12),
which includes 8 backbones within a π-stack.
19
Table 2. Summary of structural characteristics of thermally annealed thin film via GIWAXS,
and OFET device data for the isoindigo polymer series.
Thin-Film Morphology OFET Devices
Entry Polymer
Lamellar
Distance
(Å)
π-
Stacking
(Å)
Backbone
Orientation Architecture
µh µ
e
Ref
(cm2V
-1s
-1)
1 P[T-iI(OD)]-1 22, weak
Relatively
Amorphous
BG/TC,
OTS treateda
1.5×10-
2
NR 56
2 P[T-iI(HD)]-3 20 3.7 Face-on BG/TCb 1×10
-3 8×10
-3
25
3 P[T-iI(HD)]-4
20 3.8 Face-on BG/BC,
HMDS
treatedc
4×10-2
1×10-1
32
20 3.8
CL: 80Å Face-onf
4 P[T2-iI(OD)]-
1 20
Edge-on
BG/TC,
OTS treateda
0.42 NR 56
5 P[T2-iI(OD)]-
2 20
3.8
CL: 80Å Edge-on
BG/TC,
OTS treatedb
0.33 NR 22
6 P[T2-iI(OD)]-
3 20 3.8 Edge-on
BG/TC,
OTS treatedd
0.66 NR 21
7 P[T2-iI(4-
DTd)] 25 3.6 Edge-on
BG/TC,
OTS treatede
2.98 NR 21
8 P[T2-iI(SiO)] 23
3.6
CL:
138Å
Edge- and
face-on
BG/TC,
OTS treatedb
2.00 NR 22
9 P[DTS(EH)-
iI(EH)] 17 3.7 Face-on
BG/BC,
HMDS
treatedc
1×10-3
1×10-3
32
10 P[T3(C6)-
iI(HD)]-1 20 3.7 Face-on
BG/BC,
HMDS
treatedc
5×10-2
- 32
11 P[T3(C6)-
iI(HD)]-2 18
3.7
CL:
100Å
Edge-onf Results on P[T3(C6)-iI(HD)]-1 used
here
20
12 P[T3(C8)-
iI(HD)]-3 21
3.7
CL: 29Å Edge-on BG/TC
b 3×10
-2 5×10
-3
25
13
P[T3(C6)-
iI(SiO)] 25
3.6
CL: 92Å Face-on
BG/BC,
HMDS
treatedc
4×10-3 -
14 P[T6(C8)-
iI(HD)] 19 3.7 Edge-on BG/TC
b 7×10
-3 2×10
-4
25
aspin coated from chloroform, annealed at 150°C for 20 min; bspin coated from chloroform,
annealed at 170°C for one hour; cdrop cast from chloroform, annealed at 150°C for one hour;
dspin coated from tetrachloroethane, annealed at 150°C for 30 min; espin coated from
tetrachloroethane, annealed at 175°C for 30 min; fspin coated from o-dichlorobenzene, annealed
at 200°C for 5 min. All source/drain contacts are gold. BG: bottom-gate; TC: top-contact; BC:
bottom-contact; CL: coherence length for (010) peak; NR: not reported.
In developing a perspective on these results, trends can be established for this series of
copolymers based on results detailed in Table 2. A first observation is the impact of processing
conditions on the GIWAXS pattern of some polymers. Indeed films of P[T3(C6)-iI(HD)]
processed by drop-casting from a chloroform solution tend to exhibit a (100) lamellar peak along
the Qxy axis (entry 10), whereas films spin coated from oDCB exhibit a (100) peak along the Qz
axis (entry 11). By comparison, P[T-iI(HD)]-4 exhibits a (100) along the Qxy axis, regardless of
the processing conditions. As such, care needs to be made in comparing results between different
studies where structural (repeat unit, degree of polymerization, dispersity, end group identity,
branching, etc.), primary processing (solvent, temperature, substrate, time), and post-processing
(solvent and thermal annealing, use of additives, etc.) all need to be carefully accounted for.
Considering this, a second observation from these results is the lack of correlation between
21
molecular weight and packing in the case of P[T-iI(HD)] and P[T2-iI(OD)], where variations in
molecular weight do not influence the π-stacking or lamellar distances (entries 2 and 3, and 4-6).
In addition to molecular weight, polydispersity and the identity of end-capping groups are also
expected to impact the scattering results as discussed by Koldemir et al.58 However, the
challenges associated with characterization of the identity of the end-groups and the lack of
reported data on the synthetic addition of end-cappers prevents a discussion on the effect of end-
groups on polymer ordering in thin films in this article. A third observation is that interchain
packing is not influenced by the number of thiophene rings in the backbone, but rather by the
distance between the backbone and branching point on the solubilizing side-chain, with both
siloxane-containing polymers P[T2-iI(SiO)] and P[T3(C6)-iI(SiO)] exhibiting a π-stacking
distance of 3.6 Å (entries 8 and 13). Another noticeable change between the P[T3(C6)-iI(HD)]-2
and P[T3(C6)-iI(SiO)] thin film structures (Figure 3b and c) is their orientation relative to the
substrate, which was also observed in the case of P[T2-iI] by Bao et al.22, 38 Indeed, P[T3(C6)-
iI(SiO)] chains are mainly face-on ((010) π-stacking out-of-plane peak along Qz axis and (100)
lamellar scattering in-plane peak along the Qxy axis), whereas after annealing the P[T3(C6)-
iI(HD)] backbones are predominantly edge-on with regards to the substrate. It was hypothesized
by Chen et al.31 that highly soluble, planar isoindigo-like polymers lead to a majority of the
chains being face-on ((100) lamellar in-plane scattering), which can explain why the more
soluble, less sterically hindered P[T3(C6)-iI(SiO)] shows predominantly face-on orientation
compared to P[T3(C6)-iI(HD)]-2. Another explanation for the face-on orientation of P[T3(C6)-
iI(SiO)] could be the amphiphilic-like effect of the siloxane and alkyl side-chains. As illustrated
by Kim and Swager,55 it is possible that the siloxane side-chains on either side of the isoindigo
units interact preferentially with the SiO2 surface, while the alkyl side-chains on the terthiophene
22
units are elongated perpendicular to the substrate, leaving the polymer backbone to be face-on
with regards to the substrate. This initial layer of face-on polymer backbone can further act as a
template for the growth of subsequent polymer layers into a face-on arrangement. The different
orientations are expected to affect the charge transport properties of the pristine polymers.
Charge Transport. Solid state structure of thin conjugated polymer films has been shown to
have a large influence on transport properties as measured in field-effect transistors,16 but
straightforward conclusions based on structure alone are difficult to obtain due to the sensitivity
of measurements to processing conditions such as solvent effects (polymer-solvent interactions
and film solidification rate)59 and to the device structure such as substrate choice, electrode
material and contact resistance, and device architecture (bottom gate-bottom contact for e.g.).
With so many variables, care must be taken in comparing these results with the relevant
statistical analysis but general concepts can be extracted from comparing the structures described
in Table 2. To test the influence of polymer packing on transport properties in these iI-based
polymers, bottom-gate, bottom-contact OFETs were fabricated based on P[T3(C6)-iI(SiO)] and
tested, and the results are included in Table 2, which summarizes results previously reported in
the literature on similar isoindigo-based systems. In the approach used here, the active layer was
drop cast at room temperature from a solution of the polymer (2 mg/mL) in chloroform, followed
by drying at 150°C for one hour to reproduce the fabrication conditions of Stalder et al.32
Annealing of the polymer films at 150 °C and 200 °C leads to similar GIWAXS data so both
processing conditions lead to similar observations of polymer packing and conclusions as to the
relationship between film morphology and transport properties.
Previous OFET devices fabricated using these structures have shown ambipolar behavior for
P[T-iI(HD)] (entries 2 and 3), and some conflicting results for P(T3-iI)-like polymers (entries 10
23
and 12). In the case of P[T3(C6)-iI(HD)]-1, an OFET hole mobility of 5×10-2 cm2V-1s-1 was
obtained, but no n-type behavior was observed. However, when the n-hexyl chains are replaced
by n-octyl chains in P[T3(C8)-iI(HD)]-3, the electron mobility is reported to be 5×10-3 cm2V-1s-1,
with a hole mobility similar to that of P[T3(C6)-iI(HD)]-1 as seen in entries 10 and 12 of Table
2.25 Ambipolar transport was measured in some isoindigo polymers (Table 2, entries 2,3,9,12 and
14); however there is no clear correlation between electron mobility and isoindigo density in the
backbone. This is due in part to the numerous parameters which come into play when measuring
charge mobility in OFET devices. One example where device parameters are kept constant to
focus on the effects of polymer structure on transport was conducted by Stalder et al. using 2D
solid-state NMR to elucidate the interactions between polymer backbones. The authors conclude
that P[T-iI(HD)]-4 packing include slightly shifted isoindigo stacks, which maintain overlap of
the isoindigo units over multiple polymer backbones, whereas P[T3(C6)-iI(HD)]-1 packing
favors mixed interactions of isoindigo and terthiophene across polymer backbones.32 These
differences in molecular packing and orbital overlap between P[T-iI] and P[T3-iI] polymers
account in part for the ambipolar charge transport in P[T-iI] polymer, whereas P[T3-iI] tend to be
p-type polymers. This conclusion is backed up by the fact that the homopolymer of isoindigo,
prepared via Suzuki coupling, exhibits only n-type transport (η = 1×10-4 cm2V-1s-1) due the
propensity of iI interactions.32 In the case of P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-1, the
differences in charge transport can thus be explained by considering intrachain transport at the
molecular level, with breaks in conjugation observed for the LUMO of P[T3-iI], and by
interchain charge hopping through electronic coupling of adjacent iI units, as was discussed in
the case of poly(fluorene-alt-benzothiadiazole) F8BT.46, 47
24
Examining the results of Table 2, the isoindigo-based polymer films which exhibit π-stacking
have π-π distances of 3.6 Å to 3.8 Å. While a short π-stacking distance may be a necessary
condition for high interchain overlap and high mobilities, P[T-iI(HD)], P[DTS(EH)-iI(HD)], and
P[T3(C6)-iI(SiO)] (entries 2-3, 9 and 13 respectively) demonstrate that it is not a sufficient
characteristic in these polymers. Indeed, all three polymers have a π-stacking distance around 3.7
Å but two also have a face-on polymer orientation relative to the substrate, which sets the charge
hopping direction perpendicular to the transport direction in OFETs, and explains in part their
lower hole mobility compared to P[T3(C6)-iI(HD)]-1. Although orientation is not the only
determining factor in some structures, the increase in mobility with edge-on orientation in
isoindigo polymers seems to follow previous trends observed for DPP-containing polymers.16
Coherence length (as discussed in the previous paragraph, also referred to as crystallite size) in
the π-stacking direction also needs to be considered to explain the trends seen in the transport
properties. For instance, P[T2-iI(OD)] (entries 4-6) shows similar packing to P[T3(C8)-iI(HD)]-3
(entry 12), but possesses a hole mobility that is an order of magnitude higher. This is explained
by the fact that crystallinity is maintained over 3 nm (9 backbone segments) in the case of
P[T3(C8)-iI(HD)]-3 and 8 nm (22 backbone segments) in the case of P[T2-iI(OD)]-2. The
relationship between long-range order in crystallites and high mobility can be used to explain
trends seen in isoindigo polymers, as discussed broadly for conjugated polymers in general in the
recent work by Noriega et al.15 and Zhang et al..14 In one example, Zhang et al. found that both
ordered aggregates and amorphous regions of a copolymer of indacenodithiophene and
benzothiadiazole (IDT-BT) are uniformly oriented with respect to the substrate and both domains
provide pathways for efficient intra- and interchain transport.14 In another publication, Noriega et
al. showed that short-range order within polymer aggregates are connected through amorphous
25
regions by tie chains, which still allow for efficient interchain transport in the amorphous
domains.15 As such, the work reported here on isoindigo-containing polymers does give some
guidelines for the relationship between coherence length and charge carrier mobility, but our
description is limited in the sense that only the short-range ordered regions were observed by
GIWAXS and no data was obtained to include the impact of amorphous regions of the films on
charge transport.
Solar Cell Properties. To determine the impact of the ground and excited state energy levels,
along with the polymer’s solid state behavior, on photovoltaic properties, solar cells were
fabricated and the processing conditions were optimized using a conventional device structure
(indium tin oxide (ITO)/PEDOT:PSS/Isoindigo-based polymer:PC71BM/LiF/Al).
26
Table 3. Summary of optimized conventional OPV device characteristics.
Processing OPV Characteristics
Entry Polymer Poly:
PCxBM
Solvent/
Additive
Thick.
(nm)
Jsc
(mA cm-2)
Voc
(V)
FF
(%)
PCE
(%) Ref
1 P[T-iI(OD)]-2 1:2
PC71BM CB 70 1.7 0.87 60 0.9a 39
2
P[T-iI(HD)]-1
1:2
PC71BM
oDCB/
2.5%
DIO
85 5.4 0.89 63 3.0a 41
3 1:1
PC61BM
oDCB/
2.5%
DIO
80 9.1 0.91 54 4.5a 37
4 P[T-iI(HD)]-2 1:1.5
PC61BM
oDCB/
2.5%
DIO
90 5.7 0.92 63 3.3 28
5 P[T-iI(HD)]-3 1:1.5
PC61BM
oDCB/
2.5%
DIO
100 7.2 0.91 66 4.3 a 25
6 P[T-iI(HD)]-4 1:1.5
PC71BM
oDCB/
2.5%
DIO
70 5.8 0.90 55 2.9
7 P[T2-iI(OD)]-2 1:1
PC71BM
CF/
4% DIO 170 10.3 0.91 60 5.2 42
8 P[T2-iI(SiO)] 1:2
PC71BM
CF/
0.5%
DIO
150 9 0.87 52 3.4 42
9 P[DTS(EH)-iI(EH)] 1:4
PC71BM
CB/
4% DIO 8.3 0.76 42 2.6 44
10 P[T3(C6)-iI(HD)]-2 1:1.5
PC71BM
oDCB/
2.5%
DIO
100 15.2 0.70 62 6.6
11 P[T3(C8)-iI(HD)]-1 1:1.5
PC71BM
oDCB/
2.5%
DIO
90 13.1 0.70 69 6.3a 40
12 P[T3(C8)-iI(HD)]-2 1:1.5
PC71BM
oDCB/
2.5%
DIO
90 12.9 0.70 63 5.7 28
27
13 P[T3(C8)-iI(HD)]-3 1:1.5
PC71BM
oDCB/
2.5%
DIO
100 14.6 0.72 66 6.9 25
14 P[T3(C8)-iI(HD)]-4 1:1.5
PC71BM
oDCB/
3% DIO 13.8 0.73 64 6.4 26
15 P[T3(C6)-iI(SiO)] 1:1.5
PC71BM
oDCB/
2.5%
DIO
90 8.6 0.66 49 2.8
16 P[T6(C8)-iI(HD)]-1 1:1.5
PC71BM
oDCB/
2.5%
DIO
100 10.5 0.65 66 4.7 25
17 P[T6(C8)-iI(HD)]-2 1:1.5
PC71BM
oDCB/
3% CN 15.7 0.71 64 7.1 26
a Best device characteristics, all others are averages.
The optimized device results are summarized in Table 3. As can be seen, the use of 1,8-
dioodoctane (DIO) as a processing additive proves to be beneficial to the PCE in these blends. In
particular the drying time of the active layer was found to have a significant influence on device
performance. When the P[T3(C6)-iI(HD)]-2:PC71BM film was placed under vacuum
immediately after spin-coating the active layer, followed by LiF/Al deposition, the devices
showed a Voc of 0.70±0.01 V, a Jsc of 14.2±0.4 mA cm-2, and FF of 59 %. The PCE was therefore
calculated to be 6.0±0.1 %. However, after drying the active layer overnight at room temperature
in nitrogen atmosphere, the Jsc, FF, and PCE increased to 15.2±0.1 mA cm-2, 62±1.0 %, and
6.6±0.1 %, respectively. These results indicate that an increase in Jsc and FF is observed with the
long evaporation time of the oDCB:DIO mixture. If the evaporation time of the solvent is short,
there is insufficient time for self-organization of the polymer in the film as it solidifies. On the
other hand, long evaporation times allow sufficient time for self-organization to occur.60, 61
Indeed, Rogers et al.62 demonstrated little change in the assembly of poly(2,1,3-
benzothiadiazole-4,7-diyl[4,4-bis(2-ethylhexyl)-4H-cyclopenta[2,1-b:3,4-b′]-dithiophene-2,6-
28
diyl]) (PCPDTBT) polymer chains in films cast from chlorobenzene past 2 min after casting,
whereas films cast with high boiling point 1,8-octanedithiol (ODT) as a processing additive
showed structural evolution up to 78 min after casting. Furthermore, Ho et al.26 observed
increased short-circuit currents and fill factors in isoindigo-based devices, where the active layer
is processed with 1-chloronaphthalene (CN) as a processing additive. In their work, the active
layer was left to dry overnight under inert atmosphere, which could indicate further
rearrangement with longer drying times when CN is used as a processing additive. One question
that arises is the possibility of the high boiling point additive remaining in the film after casting.
The use of high vacuum (10-6 torr) on the active layer prior to evaporation of the top electrode
provides one mechanism for removal of the processing additive from the film. Previous reports
have discussed the absence of thiol groups from ODT when observed in FTIR, Raman and XPS
after drying the films under 10-3 torr.63 In this work, blends cast with DIO were analyzed by
XPS, and showed no residual iodine peak in the top nanometer layer of the bulk heterojunction
as seen in Table 5. However, complete removal of DIO from the active layer in the final device
is still not fully proven.
As observed in comparing previous reported results summarized in Table 3 to the new results
reported here, we show that the number of thiophenes in the backbone has an influence on the
device performance.25, 26 In the case of P[T-iI(HD)]-4 (entry 6) which has one backbone
thiophene, Voc is 0.9 V while the Voc of P[T3(C6)-iI(HD)]-2 (entry 10) with three backbone
thiophenes is 0.7 V. However, P[T3(C6)-iI(HD)]-2 shows higher PCEs compared to P[T-
iI(HD)]-4 due to higher Jsc and FF. The morphology investigated by atomic force microscopy as
shown in Figure 4 gives some insight as to the cause for this increase in Jsc. As seen in Figures
4b and d, P[T3(C6)-iI(HD)]-2:PC71BM exhibits finer phase separated microstructure compared
29
to P[T-iI(HD)]-4:PC71BM blends, which is one possible explanation to the higher Jsc in
P[T3(C6)-iI(HD)]-2 devices than P[T-iI(HD)]-4 devices.
Preliminary analysis of Table 3 highlights some discrepancies in Jsc trends between P[T3-iI]
device and P[T6-iI] device as shown by entries 13-14 and 16-17. In work by Ho et al.,26 CN was
used as a solvent additive instead of 1,8-diiodooctane (DIO) that was used by Ma et al.25 in
P[T6(C8)-iI(HD)]:PC71BM blend, leading to an improvement in Jsc from 10.5 mA cm-2 to 15.7
mA cm-2. As summarized in Table 3, almost all reported devices based on isoindigo-containing
polymers required the use of processing additives to achieve optimized OPV performance. As
such, the effect of processing additives on the morphology and device characteristics of P[T-
iI(HD)]-4, P[T3(C6)-iI(HD)]-2 and P[T3(C6)-iI(SiO)]-based films will be discussed in the
following paragraph.
Effect of DIO as a Processing Additive. DIO has been successfully used to control phase
separation in BHJ devices constructed using isoindigo-based polymers blended with fullerenes,
inducing a fine structure, which enhances cell performance.40, 44 As illustrated by the results in
Table 4 for the polymers studied in this work, adding DIO into isoindigo polymer:PC71BM
blends yields an enhancement in Jsc and FF, which in turn leads to higher PCE for all three
polymer systems studied. The largest increase in Jsc with DIO is observed in blends based on
P[T3(C6)-iI(SiO)], with the use of DIO leading to an over five-fold enhancement in Jsc to 8.6 mA
cm-2 compared to 1.6 mA cm-2 without DIO. The Jsc is also increased from 2.3 mA cm-2 to 5.8
mA cm-2 and from 12.5 mA cm-2 to 15.2 mA cm-2 with DIO in blends based on P[T-iI(HD)]-4
and P[T3(C6)-iI(HD)]-2 respectively. Interestingly, the Voc decreases (< 10%) for all devices
cast with DIO compared to devices fabricated without DIO. This was also observed in blends
30
containing P[DTS(EH)-iI(HD)], where the Voc dropped to 0.76 V with DIO from 0.86 V without
DIO.44
Table 4. Effect of DIO concentration in oDCB on conventional OPV device characteristics
Polymer:PC71BM Jsc
(mA cm-2)
Voc
(V)
FF
(%)
PCE
(%)
P[T-iI(HD)]-4 2.3±0.1 0.97±0.01 52±2 1.2±0.1
P[T-iI(HD)]-4, 2.5% DIO 5.8±0.1 0.90±0.01 55±1 2.9±0.1
P[T3(C6)-iI(HD)]-2 12.5±0.5 0.72±0.01 56±1 5.0±0.2
P[T3(C6)-iI(HD)]-2, 2.5% DIO 15.2±0.1 0.70±0.01 62±1 6.6 ±0.1
P[T3(C6)-iI(SiO)] 1.6±0.3 0.69±0.01 37±1 0.4 ±0.1
P[T3(C6)-iI(SiO)], 2.5% DIO 8.6±0.2 0.66±0.01 49±1 2.8±0.1
Although both P[T3(C6)-iI(HD)]-2 and P[T3(C6)-iI(SiO)] have the same conjugated backbone
structure, they express very different OPV device efficiencies.38 These differences in efficiencies
can be partly explained by the variations in lateral phase structure that cause differences in
charge generation and transport, leading to differences in Jsc and FF. Atomic force microscopy
(AFM) was used to investigate the change in surface morphology as a function of DIO content as
shown in Figure 4. It can be seen that DIO has a large effect on the topology in the case of P[T-
iI(HD)]-4, but little influence on blends with P[T3(C6)-iI(HD)]-2. The finer feature size in the
P[T-iI(HD)]-4 and P[T3(C6)-iI(SiO)] films formed with DIO are consistent with an increase in
Jsc compared to films processed without this additive. Comparing these results, it appears that
DIO as a processing additive has little visible effect on the microscopic phase separation in the
case of materials with low solubility in the processing solvent, such as P[T3(C6)-iI(HD)]-2,
when compared to the more soluble polymers P[T-iI(HD)]-4, P[T3(C6)-iI(SiO)], or P[DTS(EH)-
iI(HD)].44 However, the use of DIO does impact surface roughness in all three polymer films. To
31
a certain extent, a rougher surface of the active layer in OPV can improve light harvesting and
charge collection. Surface roughness can also be used as an indicator of crystallite formation in
thin films. The increase in surface roughness can explain in part the increase in FF and Jsc, both
leading to higher power conversion efficiency in the devices processed with DIO, although the
correlation between increased surface thickness and increased FF and Jsc remains unclear for
these three systems.
Figure 4. AFM images of blends of PC71BM with (a,b) P[T-iI(HD)]-4, (c,d) P[T3(C6)-iI(HD)]-
2, and (e,f) P[T3(C6)-iI(SiO)] without (top) and with 2.5%v DIO (bottom). The root-mean-
square roughness is indicated at the bottom right corner. All images are 1×1 µm2 with a 30 nm
height scale.
The mechanism by which processing additives modify the blend morphology is complex;64, 65
however, some preliminary work has shown that DIO has an impact on the size of PC71BM
aggregates.66 Without DIO present, fullerene derivatives tend to aggregate into large fullerene
rich domains as the film dries. In comparison, DIO serves as a better solvent for the fullerene
than chlorobenzene or oDCB, and reduces PC71BM aggregation, allowing fullerene to mix more
32
homogeneously with the polymer in the amorphous phases. In contrast to this hypothesis for
DIO, Schmidt et al. 67 have postulated that CN helps form polymer lamellar crystallites within
polymer aggregates in solution, which in turn leads to smaller and more ordered polymer
domains in the dried thin film. Gao et al.68 also demonstrated the variations in film forming
mechanisms when aliphatic additives (e.g. DIO) are used as opposed to aromatic additives (e.g.
oDCB). In the latter case, polymer aggregates with some degree of ordering are formed in
solution as observed with UV-visible absorption. Yan et al.69 have also discussed the impact of
polymer self-assembly on the choice of processing additive for optimal device performance. CN
was shown to increase Jsc compared to DIO in blends where the polymer exhibited a propensity
to form polymers fibrils, and addition of DIO leads to an increase in Jsc in blends where the
polymer forms less crystalline, irregular particles. This variation in film formation mechanism
could explain the difference in OPV device parameters in the case of P[T6(C8)-iI(HD)] devices
as measured by Ma et al.25 and Ho et al,26 where Jsc of 10.5 mA cm-2 and 15.7 mA cm-2 are
measured in blends cast with DIO and CN respectively (entries 16 and 17). In this case,
differences in self-assembly of P[T3(C8)-iI(HD)] and P[T6(C8)-iI(HD)] chains could explain the
difference in choice of processing additive between DIO and CN for optimized device
performance.
In addition to AFM, GIWAXS was conducted on blends of P[T3(C6)-iI(HD)]-2 and P[T3(C6)-
iI(SiO)] with PC71BM to check for a correlation between molecular ordering and device
parameters (Figures S4 and S5), as discussed in the previous paragraph for other polymer
systems. Based on the GIWAXS results on both polymer:PC71BM blends, slight changes are
observed in the molecular orientation and crystalline packing of the polymers in films spun cast
with and without DIO. In Figure S4, an increase in the scattering intensity is seen in both
33
polymer blends when DIO is used, suggesting an increase in crystallinity. This is also supported
by the increase in intensity of the higher order (200) peak in blends based on P[T3(C6)-iI(SiO)]
with DIO (Figures S4d and S5c). Another differences observed is a slight broadening of the
(100) peak around the Qz axis when DIO is used (Figure S5, b and d) for both polymers, As
such, the GIWAXS results on polymer crystallites do not indicate a strong dependence of the
molecular morphology of P[T3(C6)-iI(HD)]-2 or P[T3(C6)-iI(SiO)] on the presence of DIO
during processing, but small variations are seen, which can offer a partial explanation for the
increase in FF and Jsc through increased crystallinity when DIO is used.
Beyond lateral phase separation, processing additives can also affect vertical phase separation as
reported by Shao et al.70 in blends based on a copolymer of benzo[1,2-b:4,5-b’]dithiophene and
thieno[3,4-b]thiophene (PBDTTT-C-T). Depending on the vertical phase separation in the
blends, charge collection can be improved at those electrodes by minimizing leakage currents.71
However, in the case of P3HT:PCBM blends, it has been demonstrated that even an interface
made up of only 3% PCBM at the cathode enables low-resistance electron collection.72 In order
to determine the surface composition of the active layer at the PEDOT:PSS and LiF/Al
interfaces, XPS was used on the device active layer to determine the carbon, nitrogen, oxygen
and sulfur content in the top few nanometers of the bulk heterojunction. The relative atomic
content were then related back to the material concentration at the surface using the chemical
formula of the two components in order to quantify the ratio of polymer to fullerene at the
LiF/Al interface of conventional cells,73 as detailed in the supporting information. As seen in
Table 5, there is an increase of the fraction of PC71BM at the surface, near the LiF/Al cathode,
with added DIO.
34
Table 5. XPS data showing elemental composition and calculated material concentrations at the
film surface as a function of processing conditions.
P[T3(C6)-iI(HD)]-2:PC71BM C
(%)
N
(%)
O
(%)
S
(%)
Polymer
(%)
PC71BM
(%)
No DIO 92.0 2.3 2.8 2.8 83±12 16±10
2.5% DIO 92.5 1.5 3.0 3.0 69±13 30±10
The vertical phase composition was subsequently fully probed using neutron reflectivity (NR) in
order to observe a possible impact of DIO on the vertical composition profile of the active layer.
Figure 5 shows the fits to the experimental data that are used to extract the vertical scattering
length density (SLD) profile for the P[T3(C6)-iI(HD)] blends cast with and without DIO. The
SLD profiles were analyzed in terms of the vertical composition profile of P[T3(C6)-iI(HD)] and
PC71BM, which have calculated SLD values of 7.29×10-7 Å-2 and 4.3×10-6 Å-2 respectively. A
homogeneous blend of these two materials with a 1:1.5 (or 40:60) ratio has a SLD of 2.87×10 -6
Å-2 for the active layer, so the NR fits were also tested to meet this requirement to ensure
accuracy of the fit parameters. Furthermore, the NR fits shown in Figure 5b and c estimate that
the thickness of the bulk heterojunction active layer beneath the aluminum (Al) electrode is
around 125 nm and 100 nm for the blends processed without and with DIO respectively.
Profilometry on the active layer of OPV devices gave active layer thicknesses of 130±3 nm and
120±3 nm for the films processed without and with DIO respectively, which correlates well with
the values obtained from fits to the NR measurements.
In terms of vertical composition in the active layer, (Figure 5b and c) the NR fits indicate that the
active layer has a non-uniform distribution of the polymer and fullerene through the film
35
thickness. Without added DIO, the surface of the active layer is slightly PC71BM-rich, with a
composition of 52% of PC71BM, and the buried interface at the PEDOT:PSS is polymer-rich,
with a composition of 74% P[T3(C6)-iI(HD)]-2 (decreased SLD) (Table S3). When DIO is used
in the casting solution, the polymer-rich phase at the PEDOT:PSS interface becomes even more
polymer-rich (composition of the layer is 91% P[T3(C6)-iI(HD)]-2) but also becomes thinner (20
nm to a couple nanometers without and with DIO respectively). Based on these NR fits, it is
shown that DIO not only has an impact on lateral phase separation, but also on vertical
composition. However, these NR fits exhibit a different trend from XPS. XPS showed an
increase of the concentration of PC71BM at the LiF/Al interface with DIO, whereas NR fits
indicate a PC71BM-rich interface without DIO. An explanation for this discrepancy is that
capping of the active layer with the cathode material could have an effect on the composition at
that interface. Moreover, the difference in the size of the area analyzed by each technique (4 cm2
in the case of NR, and less than 0.5 mm2 in the case of XPS) could account for variations
between NR and XPS results. Indeed, as seen in Figure S3 and Table S2 of the supporting
information, there is a variation of the surface composition as the measurement is taken closer to
the electrode, where the fraction of polymer drops from 83 % to 70 % closer to the Al layer.
Figure 5. Neutron reflectivity data and fits (a) leading to scattering length density (SLD) profiles
of P[T3(C6)-iI(HD)]-2 devices spun cast without (b) and with 2.5%v DIO (c).
0.01 0.1
1E-7
1E-6
1E-5
1E-4
1E-3
0.01
0.1
1
10
R(l
og
(I))
Q (A-1)
P[T3(C6)-iI(HD)]:PC71BM, no DIO
Fit, no DIO
P[T3(C6)-iI(HD)]:PC71BM, DIO
Fit, DIO
(a)
0 500 1000 1500 2000 2500 3000
5.0x10-7
1.0x10-6
1.5x10-6
2.0x10-6
2.5x10-6
3.0x10-6
3.5x10-6
Active layer: 126 nm
Average SLD: 2.6 x 10-6 A
-2
(b)
Polymer/PC71
BM
40/60
50/50
SiOx
PEDOT:PSS
Polymer rich
Al
SL
D (
A-2)
Distance (A)
P[T3(C6)-iI(HD)]:PC71BM, no DIO
0 500 1000 1500 2000 2500 3000
5.0x10-7
1.0x10-6
1.5x10-6
2.0x10-6
2.5x10-6
3.0x10-6
3.5x10-6
Active layer: 112 nm
Average SLD: 2.7 x 10-6 A
-2
Polymer/PC71
BM
40/60
Polymer rich
SiOx
SL
D (
A-2)
Distance (A)
P[T3(C6)-iI(HD)]:PC71BM, DIO
Al
PEDOT:PSS
50/50
(c)
36
Balance between Voc and Jsc. Interestingly, devices based on P[T-iI(HD)]-4 exhibit high open-
circuit voltages, but low short-circuit currents, while the opposite is true for P[T3(C6)-iI(HD)]-2.
In general, as described in Table 3, high Voc values of up to 0.91 V are obtained with blends
containing less electron-rich polymer donors (Table 3, entries 1-8), where the high Voc without
sacrificing Jsc leads to 4.5% power conversion efficiency (Table 3, entry 3).25 On the other hand,
in a device containing more electron-rich isoindigo polymer (Table 3, entries 9-17) the Voc was
lowered to 0.7 V while Jsc reached 15.7 mA cm-2, yielding higher PCEs of 7.1 % when CN is
used as a processing additive (Table 3, entry 17).26
As described in previous work,28 device results show a substantial 0.2 V difference in Voc
between P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2 or P[T3(C8)-iI(HD)] devices. Because of the
unstable electrochemical oxidation of the polymers, the understanding of this difference in Voc
based solely on the ionization potential of the polymers is challenging. Moreover, the UV-PES
results show a 0.1 eV difference in the IP of P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2, which does
not explain the 0.25 V difference in the Voc of blends cast without DIO. Vandewal et al.28 have
discussed this difference in the Voc by measuring the energy difference between the excited state
of the donor and the charge transfer state (CT) by Fourier-transform photoelectron spectroscopy
and electroluminescence.28 They find that P[T-iI(HD)]-2:PC61BM blends exhibit a lower energy
offset between the polymer excited state and the CT state (close to 0 eV) compared to P[T3(C8)-
iI(HD)]:PC71BM blends, along with low radiative charge transfer state recombination. This leads
to low voltage losses in P[T-iI(HD)]-2 cells, resulting in a high Voc of 0.9V in the case of P[T-
iI(HD)]:PC61BM blends. Furthermore, field-dependent photoluminescence quenching indicates
that this low energy difference between excited and CT states in P[T-iI(HD)]-2:PC61BM blends
37
is responsible for the low Jsc due to repopulation of the exciton from the CT state, rather than
charge separation.
In this current work, charge modulated electroabsorption spectroscopy (CMEAS) is used to
understand the change in Voc between P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2. CMEAS is a
technique that detects the changes in optical absorption below the polymer bandgap by
modulating the electric field to directly measure the effective band gap in a bulk heterojunction
solar cell.74 This novel technique allows detection of the CT states which have been correlated
with Voc. CMEAS provides higher signal-to-noise resolution compared to linear optical
absorption techniques, and takes into account interfacial effects between the two compounds in
contrast to electrochemical measurements.
The results obtained by Lai et al. using CMEAS show that the effective energy gap in P[T-
iI(HD)]-4:PC71BM blends is measured to be 1.3 eV compared to 1.0 eV in P[T3(C6)-iI(HD)]-
2:PC71BM blends.49 The 0.3 eV difference between these two blends is in agreement with the
variations observed for the Voc of the OPV devices. The CMEAS results indicate that, while
polymers may have similar onsets of oxidation, the CT states formed at the donor/acceptor
interface can have different energy levels. This difference in CT energy can be explained by the
differences in electronic coupling at the interface between polymer and fullerene in blends of
P[T-iI(HD)]-4:PC71BM and P[T3(C6)-iI(HD)]-2:PC71BM respectively, as described for other
systems by Graham et al,75 which affects the effective energy gap. The difference between CT
state energy and Voc is further explained by energy losses to reach charge separated states from
the CT state. Furthermore, higher dissociation efficiency is observed as a result of the more
delocalized CT states in P[T3(C6)-iI(HD)]-2:PC71BM blends than P[T-iI(HD)]-4:PC71BM
blends, resulting in a higher Jsc in solar cells.76 This underlines the role played by dielectric
38
constants in isoindigo-based blends, which could also explain why the low energy offset between
excited and CT states still leads to high internal quantum efficiencies (IQE) in the case of
P[T3(C6)-iI(HD)] or P[T3(C8)-iI(HD)].
Inverted Device Architecture. The effect of device architecture on device properties was
further investigated in inverted devices, more amenable to roll-to-roll fabrication. Inverted bulk-
heterojunction solar cells with a structure of ITO glass/ZnO–PVP nanocomposite/isoindigo-
based polymer:PC71BM /MoOx/Ag was fabricated using the three focus polymers in this work:
P[T-iI(HD)]-4, P[T3(C6)-iI(HD)]-2, and P[T3(C6)-iI(SiO)]. We previously reported ZnO–
poly(vinyl pyrrolidone) (PVP) composite sol–gel film5 as the ETL. On top of the ZnO–PVP
composites, the active layer in inverted devices was deposited using the same conditions as in
conventional devices. The inverted device results are summarized in Figure 6 and Table 6. For
both P[T-iI(HD)]-4 and P[T3(C6)-iI(HD)]-2, the conventional and inverted devices showed
similar device performances. On the other hand, inverted devices based on P[T3(C6)-iI(SiO)]
showed higher device performance, such as higher Jsc and FF in spite of having a lower Voc,
compared to conventional devices. The increase in device efficiency in P[T3(C6)-
iI(SiO)]:PC71BM inverted OPV devices could not only result from a difference in electrode work
function, but also from variations in the blend morphology and carrier recombination
dynamics.77
39
Figure 6. J-V curves of conventional and inverted solar cells based on P[T-iI(HD)]-4, P[T3(C6)-
iI(HD)]-2, and P[T3(C6)-iI(SiO)] polymer:PC71BM blends.
Table 6. Conventional and inverted device characteristics.
Device Jsc
(mA cm-2)
Voc
(V)
FF
(%)
PCE
(%)
P[T-iI(HD)]-4 conventional 5.8±0.1 0.90±0.01 55±1 2.9±0.1
P[T-iI(HD)]-4 inverted 5.6±0.1 0.81±0.01 52±1 2.4±0.1
P[T3(C6)-iI(HD)]-2 conventional 15.2±0.1 0.70±0.01 62±1 6.6±0.1
P[T3(C6)-iI(HD)]-2 inverted 15.2±0.3 0.68±0.01 63±1 6.5±0.2
P[T3(C6)-iI(SiO)] conventional 8.6±0.2 0.66±0.01 49±1 2.8±0.1
P[T3(C6)-iI(SiO)] inverted 9.5±0.1 0.64±0.01 63±1 3.8±0.1
SUMMARY AND PERSPECTIVE
-0.2 0.0 0.2 0.4 0.6 0.8 1.0
-16
-12
-8
-4
0
4
8
12
16C
urr
en
t d
en
sit
y (
mA
/cm
2)
Voltage (V)
P[T-iI(HD)], Conventional
P[T-iI(HD)], Inverted
P[T3(C6)-iI(HD)], Conventional
P[T3(C6)-iI(HD)], Inverted
P[T3(C6)-iI(SiO)], Conventional
P[T3(C6)-iI(SiO)], Inverted
40
In summary, we have synthesized and characterized a series of new isoindigo polymers to
highlight structure-processing-property relationships which can lead others in further developing
this family of polymers. Facile access to isoindigo in bulk syntheses suggests these systems are
worth optimization and bulk preparation for further practical development of OPV’s. These
polymers were designed to include one or three thiophene units as electron-rich moieties, and to
study the impact of side-chain branching on the electronic properties of the resulting polymer
and on the solid-state morphology. In terms of electronic and geometric considerations, the
strong electron accepting character of isoindigo localizes the electron density in the LUMO and
sets the electron affinity around 3.9 eV. The impact of side-chains on these interchain
interactions was discussed using UV-Vis absorption, 2D NMR and GIWAXS data, where
increasing the density of branched side-chains along the backbone tends to decrease interchain
interactions. These interchain interactions are particularly important for reaching high mobility
values in OFETs, where charge transport was shown to be consistently improved by increasing
the order of the isoindigo-containing polymer in thin-films, as quantified here by the coherence
length. This can be achieved by promoting backbone interactions through reducing steric
hindrance by moving the side chain branching point away from the backbone as shown in this
article by the P[T2-iI] family, and in other literature reports.78, 79 Orientation of polymer chains in
the bulk of thin films can be dictated by the chemical nature and density of the side chains, with
alkylated side chains mainly leading to edge-on orientations compared to siloxane side chains,
which yield some face-on orientation. This leads to increased transport in OFETs when the iI-
containing polymer backbones are preferentially oriented perpendicular to the substrate, and the
charge carrier mobility depends on the coherence length rather than on the π-stacking distance in
these isoindigo-based systems. Furthermore, the interchain interactions are crucial to the type of
41
charge carrier transport and to the magnitude of the charge carrier mobility, where 2D NMR was
used to correlate the iI-iI interchain interactions in P[T-iI] polymers to ambipolar charge
transport, whereas T3-iI interactions were prevalent in P[T3-iI] polymers leading to p-type
charge transport.
In OPV devices, high Voc can be achieved thanks to high ionization potentials of the isoindigo-
based polymers, which needs to be balanced with a high density of charge transfer states to
achieve high Jsc, and thus optimal PCE. The morphology in isoindigo-based blends can be tuned
by using solvent additives such as DIO or CN; however, in the case of P[T3(C6)-iI(HD)]-2 the
use of DIO does not lead to a large change in phase separation as is the case for P[T-iI(HD)]-4 or
P[T3(C6)-iI(SiO)] as investigated by AFM. The increased solubility of P[T-iI(HD)]-4 and
P[T3(C6)-iI(SiO)] in oDCB compared to P[T3(C6)-iI(HD)] is one explanation for the more
visible change in morphology when a poor solvent for the polymer, such as DIO, is used. These
observations point to the impact of the state of the polymer chains in solution prior to film
formation on the resulting thin-film morphology, in addition to the impact of residual of high
boiling point solvents in the film to allow for polymer chain rearrangement or fullerene
aggregation. This work, along with previous reports,25, 26, 40 also highlights the balance in OPV
performance, with PCEs over 6% in both conventional and inverted devices with iI-based
polymers, and the synthetic accessibility of isoindigo-containing polymers.
Overall, this work establishes some structure-processing-property relationships that could lead to
the design of novel high performance isoindigo-based polymers. Furthermore, this work can be
extended to other dually substituted acceptors such as DPP or TiI, where most of the differences
between these three units stem from their electronic characteristics, where iI-based polymers
tend to exhibit the highest ionization potential and electron affinity.
42
ASSOCIATED CONTENT
Supporting Information. Experimental details, electrochemical details and square-wave
voltamograms, description, OFET and OPV device fabrication, XPS sample preparation and
analysis, additional data from DFT calculation, monomer and polymer synthesis and gel
permeation chromatograms are supplied as Supporting Information. This material is available
free of charge via the Internet at http://pubs.acs.org.
AUTHOR INFORMATION
Corresponding Author
* Email: [email protected]
Present Addresses
† Current address: Department of Materials Science and Engineering, North Carolina State
University, Raleigh, NC, USA
Author Contributions
The manuscript was written through contributions of all authors. All authors have given approval
to the final version of the manuscript.
Funding Sources
JRR acknowledges funding of this work from the Office of Naval Research (N00014-14-1-
0173).
Notes
The authors declare no competing financial interest.
43
ACKNOWLEDGMENT
JRR acknowledges funding of this work from the Office of Naval Research (N00014-14-1-
0173). CG acknowledges the Center for Organic Photonics and Electronics for a fellowship and
R. S. acknowledges the University Alumni Awards Program for a fellowship.
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56
For Table of Contents use only:
Structure-Property Relationships Directing Transport and Charge Separation in Isoindigo
Polymers
Caroline Grand, Sujin Baek, Tzung-Han Lai, Nabankur Deb, Wojciech Zajaczkowski, Romain
Stalder, Klaus Müllen, Wojciech Pisula, David G. Bucknall, Franky So, John R. Reynolds